Silicon Carbide Materials Processing and Applications in Electronic Devices Part 3 pdf

Metastable Solvent Epitaxy of SiC,
the Other Diamond Synthetics 7
liquid
Ni
diamond
graphite
ba
temperature [K]
atomic fraction of carbon
1800
1750
1700
1650
Ni
0 0.1 0.2 0.996 0.998 1.0
graphite
graphite + diamond
1728 K
diamond
1740 K
liquid + diamond
liquid Ni(C)
α
α
+
liquid
1667 K (Ni + diamond)
1661 K (Ni + graphite)
liquid +
graphite
Fig. 6. Schematic drawings of a diamond synthesis and b enlarged sections of the Ni-C phase

diagram at 54 Kbar. In a, graphite, liquid Ni and diamond are the source, solvent and seed,
respectively. In b, the stable Ni-diamond and metastable Ni-graphite reactions are
represented by solid and dashed lines, respectively(Strong & Hanneman, 1967).
is driven spontaneously by the diffusion between small (S) and large (L) precipitates due to
Gibbs-Thomson e f fect. In thi s section, the matrix and precipitate p hases are represented as α
and β, respectively. The precipitates show the same composition of x
β
but different pressures
of small and large radii, which makes the different free energies, G
s
β
and G
L
β
,asshownin
free-energy vs. concentration diagram of Fig. 7a. The common tangents of free-energy curves
between matrix solution and precipitates show different angles and touch at the different
concentrations x
L
α
and x
S
α
of the free energy curve of the matrix solution. This concentration
difference of matrix appears at the interfaces o f the precipitates. Fig. 7b shows the schematic
diagram of the solute concentration profile of the s ystem o f α matrix and small(S) and large(L)
β precipitates. The matrix concentration equilibrating with the small precipitates should be
higher than that with the large precipitates. This difference drives the solute diffusion and
thus the simultaneous growth and dissolution of precipitates.
3.4 Direct mel ting of metastable phases

Ostwald ripening is the reaction in a solid solutions, which means the life time of metastable
phase may be longer at lower temperatures and l ower d iffusivity. Does such a m etastable
phase directly melt in liquid?
The typical example of the behavior of the coexistence of stable and metastable phases is
observed in the Fe-C system. As mentioned before, the double phase diagram of Fe-C and
Fe-Fe
3
C systems is well studied and established. The m elting behavior of metastable Fe
3
C
phase has investigated in detail by Okada et al. (1981). They measured the differential thermal
analysis (DTA) curves for the white, gray and mixture cast irons at the eutectic temperature
and composition region. Fig. 8 shows the summarized results of DTA curves as well as the
schematic double phase diagram. The endothermic temperatures shift due to the kinetic
reason of the measuring apparatus, but the corrected temperatures show the stable and
metastable eutectic temperatures o f 1426K and 1416K, respectively. The specimens of gray
cast iron contains stable phases of graphite and fcc-Fe(autenite), where they all melt only
59
Metastable Solvent Epitaxy of SiC, the Other Diamond Synthetics
8 Will-be-set-by-IN-TECH
Concentration
Free energy G
x
α
L
x
α
S
x
β

x
β
x
α
L
x
α
S
G
β
L
G
β
S
Concentration
Distance
a
b
l
G
α
Fig. 7. Schematic drawings of a free energy vs. concentration diagram and bconcentration
profile of the Ostwald ripening.
at the stable eutectic te mperature of 1426K. On the other hand, the specimens of white cas t
iron shows the double peaks of endothermic reactions at the slow heating rates. Okada et al.
(1981) found that at the first peak the metastable Fe
3
C melts but soon graphite solidifies and
then remelt at the second p eak. At the faster heating rates, o nly the melting of the metastable
Fe

3
C phase occurs. For the specimens of mixture cast iron, the reactions are complicated but
the melting and s olidifying occur simultaneously. These experimental results i ndicate that the
metastable phase is so stable that can melt directly.
Temperature
Exothermic Endothermic
gray cast iron
white cast iron
mixture iron
ΔT
a
b
Fig. 8. a DTA curves of cast irons(Okada et al., 1981) and b the double phase diagram of
equiblibrium Fe-Graphite and metastable Fe-Fe
3
Csystems.
3.5 Speculated mechanism
From the experimental result shown in Fig. 4, 4H-SiC is expected to be more stable than
3C-SiC. The Si-C system should show a double phase diagram, as schematically shown in
Fig. 9a. The corresponding free-energy vs. concentration diagram is also illustrated in Fig. 9b.
60
Silicon Carbide – Materials, Processing and Applications in Electronic Devices
Metastable Solvent Epitaxy of SiC,
the Other Diamond Synthetics 9
The solubility limit of each phase i s determined by the tangent common to the free-energy
curves of the coexisting phases. At the temperature indicated by the dotted line, the solubility
limit of metastable 3C-SiC is x
3C
l
, which corresponds to the dashed line of the liquidus in

Fig. 9a. The solubility limit of stable 4H-SiC is x
4H
l
, which corresponds to the solid line of the
liquidus. The concentration gradient in the layers consisting of 3C-SiC and 4H-SiC and a very
thin layer of liquid Si in between is obtained. The schematic carbon concentration profile in
the liquid Si layer is shown in Fig. 9c. The concentration gradient at the metastable solubility
limit of 3C -SiC leads to the extraction of C from the source plates. This concentration gradient
also causes carbon diffusion across the liquid Si s olvent and to the interface of the seed, where
C is deposited due to the supersaturation of 4H-SiC. Although the small solubility limit of C
in liquid Si, which is the cause of the slow growth of SiC in the conventional liquid method,
still remains, the small thickness of the Si solvent in the new method leads to a sufficient
concentration gradient for the growth of 4H-SiC crystals.
4. Phase stability of SiC polytypes
4.1 Phase diagram assessment of Si-C system
The key data in order to rationalize this novel process is the phase stability or the hierarchy of
SiC polytypes. The reported phase diagrams, however, are somewhat conflicting.
The standard data book Chase (1998) shows the standard formation enthalpies for α and β
phases, as follows:
Δ
f
H(α −SiC, 298.15K)=−71.546 ±6.3kJ/mol
Δ
f
H(β −SiC, 298.15K)=−73.220 ± 6.3kJ/mol.
Although the data shows that the β(cubic) phase is more stable than the α(hexagonal) phase
at 298.15K, the difference of the measured values are within the measurement errors. The
α(hexagonal) phase indicated in Chase (1998) is 6H, but also mentioned that the many
polytypes have not been adequately differentiated thermodynamically. The heat capacity and
Gibbs free energy are also reported as shown in Fig. 10. The measured values and the adapted

functions in Chase (1998) suggest that α(hexagonal) phase is less stable up to 2000K, and they
concluded unlikely the transformation to β(cubic) phase at abo ut 2300K.
The most widely adapted phase diagram should be that by Olesinski & Abbaschian (1996)
as shown in Fig. 1, where the β(cubic) phase is more stable than the α(hexagonal) phase
at any temperatures below the periodic temperature of the decomposition of SiC, 2545
◦
C.
Although the evaluators of Olesinski & Abbaschian (1996) mentioned nothing on the types of
α(hexagonal) phase, the same authors reported the co-existence o f polytypes of α phases, 6H,
15R, and 4H(Olesinski & Abbaschian, 1984). Furthermore, it also mentioned on the report
of Verma & Krishna (1966), the existence of α stability above 2000
◦
C. On the other hand,
Fromm & Gebhardt (1976) reported the different type of phase diagram as shown in Fig. 11,
wherethephasetransitionfromβ to α phases occurs at around 2000
◦
C.
Solubilities of carbon in liquid silicon measured by Hall (1958), Scace & Slack (1959), Dash
(1958), Dolloff (1960), Nozaki et al. (1970), Oden & McCune (1987), Suhara et al. (1989),
Kleykamp & Schumacher (1993), Iguchi & Narushima ( 1993), O ttem (1993), and Yanabe et al.
(1997) are summarized as in Figs. 12. Tw o reported phase diagrams as shown i n Fig. 1 and
Fig. 11 are based on the data given by Dolloff (1960). Dolloff (1960)’s data, however, are
distinctively different from the others, where the solubility limits are larger than the others.
61
Metastable Solvent Epitaxy of SiC, the Other Diamond Synthetics
10 Will-be-set-by-IN-TECH
temperature
liquid Si
SiC
a

b
free energy
liquid Si
3C-SiC
4H-SiC
μ
C
4H
μ
C
3C
x
l
3C
x
l
4H
distance
3C-SiC source
polycrystals
4H-SiC
fine particle
liquid Si solvent
carbon concentration
c
carbon concentration
carbon concentration
Si
S
+ SiC

4H
Si
S
+ SiC
3C
Fig. 9. Schematic drawings of a the p redicted Si-C double phase diagram, b related
free-energy vs. concentration diagram and c carbon concentration p r ofile in the liquid Si
solvent between the 3C-SiC source and the 4H-SiC fine particles. The metastable eutectic
temperature of the reaction liquid Si
→ Si
S
+ SiC
3C
is lower than the stable eutectic
temperature of the reaction liquid Si
→ Si
S
+ SiC
4H
,whereSi
S
denotes solid Si. The chemical
potentials of C, μ
c
, are given by the intersections of the co mmon tangents with the
pure-carbon line in b, and are spatially different in the liquid Si solvent contacting with
3C-SiC and 4H-SiC in c. The configuration of c is re lated to that of the panel on the left-hand
side in Fig. 2b.
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Silicon Carbide – Materials, Processing and Applications in Electronic Devices

Metastable Solvent Epitaxy of SiC,
the Other Diamond Synthetics 11
Fig. 10. Heat capacity and Gibbs energy of SiC(Chase, 1998).
C [at%]
L
3600
3200
2800
2400
2000
1600
1200
1414°C
2830°C
Temperature [°C]
020406080100
L+C
L+α-SiC
L+β-SiC
α-SiC+C
β-SiC+C
Si(s)+β-SiC
Fig. 11. Phase diagram of Si-C binary system including the phase transition from β to α phase
around 2000
◦
C(Fromm & Gebhardt, 1976).
Fig. 12. Solubility of carbon in liquid silicon.
63
Metastable Solvent Epitaxy of SiC, the Other Diamond Synthetics
12 Will-be-set-by-IN-TECH

The reported phase stability between α(hexagonal) and β(cubic) might be 6H and 3C. If this
assumption is true, 4H stability has not been shown experimentally. Furthermore, the distinct
difference of solubility limits indicates that the coexistence of stable and metastable phases.
4.2 Difficulty of the equilibrium state
Although the co nflicts of phase diagrams shown above remain, there have been many
attempts of experimental and theoretical researches on the kinetic p rocess of crystal growth
of SiC polytypes. Famous stability diagrams of SiC polytypes proposed by Knippenberg
(1963) and Inomata e t al. (1968) show that the crystalized phases are controlled both by the
temperature and growth rate of the operations.
The l imit to slow growth rate of kinetic processes or results of long period holding should be
equal to static results. But it is very difficult to dissolve w hole amount of SiC crystals due to
small solubility limit of SiC in liquid Si. If there remain seeds of metastable phases during the
previous pr ocesses, it is difficult to re move all of the m. T he metastable phase also grows due
to the co-exsitence of less stable phases as shown in Ostwald ripening, or from super-saturated
liquid Si. Furthermore, the required high temperature and inert environment make the static
conditions very difficult.
Inomata et al. (1969) performed careful experimental observations on the relationship
between the polytypes of SiC and the supersaturation of the solution at 1800
◦
Cwiththe
solution method, and have shown the following results;
1. β-SiC crystallizes from highly supersaturated solution. The crystals obtained at the
condition of low supersaturation, however, consist of mainly α-SiC such as 4H, 15R and
6H.
2. Relative amount of 4H increased markedly with decreasing the supersaturation.
3. From the results stated above, it is concluded that 4H is the most stable structure at 1800
◦
C
among the basic polytypes of SiC, 3C, 4H, 15R and 6H.
Those results indicate that the difference between 4H and 6H is crucial for determining the

hierarchy of SiC polytypes.
Izhevskyi et al. (2000) summarized not only the kinetic observations, but also pointed out
the impurity effects, especially nitrogen affects the transformations among 6H, 3C and 4H
SiCs. Not only through the contamination of the higher temperature operations, but also from
the starting materials made by Acheson method, specimens contain non-negligible nitrogen.
Very recent improvements on materials and apparatuses m ake it possible to avoid nitrogen
inclusions and get the hierarchy of pure SiC polytypes experimentally soon.
4.3 First principles calculations
For some cases of hardly measuring experimental value, the first principles calculations
give some hints of the puzzles, and have been applied on the topic of the hierarchy of SiC
polytypes. Liu & Ni (2005) have summarized the results o f the first principles cal culations of
SiC. All calculations show that β-SiC is less stable than α-SiCs of 2H, 4H and 6H. The hierarchy
between 4H and 6H is subtle; two of nine calculations shows that 6H SiC is most stable, but the
majority of the results indicates that 4H SiC is most stable. Of course the calculating re sults
should be judged by the precisions, the energy differences, however, are too small from 0.2
to 2 meV/Si-C pair to identify. Although the reliability of the first principles calculations
of the hierarchy of SiC polytypes are insufficient, it is important that the cal culating re sults
show against the experimental results. Those are ground state results, which means that i s
64
Silicon Carbide – Materials, Processing and Applications in Electronic Devices
Metastable Solvent Epitaxy of SiC,
the Other Diamond Synthetics 13
only reliable at low temperatures. For including the finite temperature effect, the vibrational
entropy effect is cruci al, because the configurational en tropies should be s imilar among these
SiC polytypes due to the similar local configurations of tetrahedrons. One calculating result
including the vibrational effect is shown by Nishitani et al. (2009) as in Fig. 13.
-0.002
0
0.002
0.004

0.006
4H
6H
3C
2H
Temperature [K]
Free energy difference [eV/SiC-pair]
0
1,000
2,000 3,000
Temperature [K]
Fig. 13. First principles calculations of temperature dependency of free energy difference of
6H, 3C, and 2H against 4H Si C(Nishitani et al., 2009). Finite temperature effects are included
through the vibrational free energy calculated by Phonon codes(Medea-phonon, n.d.;
Parlinski et al., 1997).
These first principles calculations were carried out using the Vienna Ab initio Simulation
Package (VASP) code(Kresse & Furthmüller, 1996a;b; Kresse & Hafner, 1993; 1994). The
interaction between the ions and valence electrons was described by a projector
augmented-wave (PAW) method(Kresse & Joubert, 1999). A plane-wave basis set with a cutoff
of 400 eV was used. The exchange-correlation functional was described by the generalized
gradient approximation (GGA) of the Perdew-Wang91 form(Perdew & Wang, 1992). Phonon
calculation was performed by a commercial pre-processor of Medea-phonon(Medea-phonon,
n.d.) with t he direct method developed by Parlinski et al. ( 1997). The volumes and/or c/a
ratios were fitted to the most stable point at each temperature.
Fig. 13 shows the te mperature dependencies of free energy of 6H, 3C, and 2H SiC polytypes
measured from 4H SiC. 4H SiC is most stable at low temperatures, but 6H Si C is most stable
at higher temperatures. 3C SiC is less stable against 4H or 6H SiC except at very high
temperature region. Those results are consistent with the other speculations but the precisions
of the calculations are not enough. Although the more precise calculations will alter the results
of hierarchy of polytypes, their result pointed out the possibility of the phase transition in the

Si-C system from the first principles calculations.
5. Conclusions
We have utilized a new method for manufacturing SiC from liquid Si; in this method,
single crystals of 4H-SiC are obtained from polycrystalline 3C-SiC source in the absence
of a temperature gradient. The origin of the driving force for crystal growth is the same
as that in the case of diamond synthesis from a metal-carbon solvent, and it is elucidated
by considering the stable-metastable double phase diagrams. This similarity in the growth
mechanism indicates that the methods developed for diamond synthesis can be directly used
for growing large-size SiC crystals from a metastable solvent of Si.
65
Metastable Solvent Epitaxy of SiC, the Other Diamond Synthetics
14 Will-be-set-by-IN-TECH
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the Other Diamond Synthetics 15
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Silicon Carbide – Materials, Processing and Applications in Electronic Devices
4
The Formation of Silicon Carbide in the SiC
x

Layers (x = 0.03–1.4) Formed by Multiple
Implantation of C Ions in Si
Kair Kh. Nussupov and Nurzhan B. Beisenkhanov
Kazakh-British Technical University
Kazakhstan
1. Introduction
Promising application of thin-film technology is the synthesis of SiC, possessing such
valuable properties as high hardness (33400 Mn/m
2
), chemical resistance, high melting
point (2830°C), wide bandgap (2.3–3.3 eV), etc (Lindner , 2003). Unfortunately, since it is still
difficult to grow SiC material of crystalline quality to meet requirements for a large scale
industrial application, small-size and high-cost SiC wafers severely limit their applications
at present (Liangdeng et al., 2008). Doped with different impurities, silicon carbide is used in
semiconductor technology (Yаn et al., 2000; Chen et al., 2003). Field-effect transistors, diodes
and other electronic devices based on SiC have several advantages compared to similar
silicon devices. Among them, the opportunity to work at temperatures up to 600°C, high
speed and high radiation resistance. A large number of polytypes of SiC makes it possible to
create heteropolytype structures (Lebedev et al., 2001, 2002a, 2005) to form a defect-free,
near-perfect contacts with unusual electronic properties (Fissel et al., 2001; Lebedev et al.,
2002b; Semenov et al., 2010). Diode structures have been established (Lebedev et al., 2002b),
in which the value of uncompensated donors N
d
−N
a

was (1.7−2)×10
17
см
-3
in the layer (n)
6H-SiC and acceptors N
a
−N
d
~ 3×10
18
см
-3
in the layer (p) 3C-SiC. In the spectrum of the
electroluminescence of diodes revealed two bands with maxima hν
max
≈ 2.9 eV (430 nm) and
2.3 eV (540 nm), close the band gaps of 6H-and 3C-SiC. Currently, using the methods of
vacuum sublimation (Savkina et al., 2000), molecular beam epitaxy (Fissel et al., 1996), the
epitaxial and heteropolytype layers based on the cubic 3C-SiC and two hexagonal 6H-SiC,
4H-SiC on substrates of SiC, are grown. By chemical vapor deposition (CVD) (Nishino et al.,
2002) are grown heteroepitaxial layers of 3C-SiC on substrates of Si. At the temperatures
below 1200°C there are conditions for the growth of both poly- and nanocrystalline SiC with
different degrees of crystallinity and structure of the cubic polytype 3C-SiC. Such conditions
were realized in the magnetron sputtering (Kerdiles et al., 2000; Sun et al., 1998), laser
ablation (Spillman et al., 2000) and plasma deposition (Liao et al., 2005), plasma-enhanced
chemical vapor deposition (George et al., 2002; Pajagopalan et al., 2005), molecular beam
epitaxy (Fissel et al., 2000). At temperatures below 1500°C in the direct deposition of carbon
and silicon ions with an energy of ~100 eV, the growth of nanocrystalline films with a
consistent set of the polytypes 3C, 21R, 27R, 51R, 6H is possible (Semenov et al., 2008, 2009,

2010). Photoluminescence spectrum from the front surface of the nanocrystalline film

Silicon Carbide – Materials, Processing and Applications in Electronic Devices

70
containing cubic 3C and rhombohedral 21R, has a band emission with three peaks at 2.65,
2.83, 2.997 eV (469, 439, 415 nm), the shoulder of 2.43 eV (511 nm) and a weak peak at 3.366
eV (369 nm) (Semenov et al., 2010).
Of particular interest is the synthesis of SiC layers in silicon by ion implantation (Lindner,
2003; Liangdeng et al., 2008; Kimura et al., 1982, 1984) due to the possibilities to obtain the
films of a given thickness and composition, nanolayers of chemical compounds and
multilayer structures, as well as the insulating layers in the manufacture of integrated
circuits. Silicon structures with a hidden layer of silicon carbide can be used as a SOI-
structure (silicon-on-insulator), which have advantages over structures with a hidden layer
of SiO
2
, obtained by the SIMOX (Separation by IMplantation of OXygen). Crystalline films
of high-quality β-SiC on SiO
2
can be obtained by multiple ion implantation of C in Si and
selective oxidation of the top layer of Si (Serre et al., 1999).
High-dose carbon implantation into silicon in combination with subsequent or in situ
thermal annealing has been shown to be able to form polycrystalline or epitaxial cubic SiC
(β-SiC) layers in silicon (Liangdeng et al., 2008; Kantor et al., 1997; Durupt et al., 1980;
Calcagno et al., 1996). To create a SiC−Si heterojunctions were among the first used the
method of ion implantation the authors of (Borders et al., 1971; Baranova et al., 1971), in
which the synthesis of silicon carbide was carried out by implantation of
12
C
+

ions (Е = 200
кэВ, D = ~10
17
см
-2
(Borders et al., 1971) or Е = 40 кэВ, D > 10
17
ион/см
2
(Baranova et al.,
1971)) into a silicon substrate. From the position of the IR absorption band (700–725 cm
-1
),
reducing its half-width after annealing and displacement in the region at 800 cm
-1
,
corresponding to transverse optical phonons of SiC, it was found that the formation of
crystalline SiC phase occurs in the temperature range near 850ºC (Borders et al., 1971) and
900°C (Baranova et al., 1971). Silicon carbide was identified using transverse optical phonon
spectra in most of the above work on ion implantation, as well as in (Gerasimenko et al.,
1974, Wong et al., 1998, Akimchenko et al., 1977a, 1980; Chen et al., 1999, Kimura et al.,
1981). The detection of longitudinal optical vibrations of lattice atoms (LO phonons) and
their changes during film annealing give additional information on the crystallization
processes (Akimchenko et al., 1977b, 1979).
Difficulties associated with the problem of synthesis of a crystalline silicon carbide
prevent the wide use of SiC in microelectronics. The Si-C mixture, after implantation of
large doses of carbon, is assumed to be amorphous (Lindner , 2003; Liangdeng et al., 2008;
Kimura et al., 1981). Carbon atom diffusion in the implanted layer is restricted by the
strong Si-C bonds (Liangdeng et al., 2008). A negative influence of stable C- and C−Si-
clusters (Yan et al., 2000, Chen et al., 2003; Kimura et al., 1982, 1984; Durupt et al., 1980,

Calcagno et al., 1996; Borders et al., 1971) on the crystallization of SiC in the implanted
silicon layers with different concentration of the implanted carbon, was found. Heat
treatment up to 1200°C does not lead to complete disintegration of the clusters and the
release of C and Si atoms to form the Si–C-bonds with tetrahedral orientation which is
characteristic of the crystalline SiC phase (Khokhlov et al., 1987). However intensively
developing area is the formation in SiO
2
by ion implantation of nanostructured systems
with inclusions of nanocrystals and clusters of Si, SiC and C, providing the expense of size
effects luminescence virtually the entire visible spectrum (Zhao et al., 1998; Tetelbaum et
al., 2009; Perez-Rodrıguez et al., 2003; Gonzalez-Varona et al., 2000; Belov et al., 2010).
This makes it necessary to study the mechanisms of formation of nanocrystals Si, SiC,
carbon nanoclusters and amorphous SiC precipitates during the implantation and
The Formation of Silicon Carbide in the SiC
x
Layers
(x = 0.03–1.4) Formed by Multiple Implantation of C Ions in Si

71
annealing. It is of considerable interest to study the effect of the concentration component,
nanoclusters, the phase composition of SiC films and their heat or plasma treatment on
the crystallization processes and clustering, the size of nanocrystals and, consequently,
the physical properties of the films.
A number of studies have shown that the implantation of ions with multiple different
energies is the most suitable for obtaining a uniform layer of carbon and silicon atoms to
form SiC (Liangdeng et al., 2008, Calcagno et al ., 1996, Srikanth et al., 1988,
Rothemund&Fritzsche, 1974, Reeson et al., 1990, Lindner et al., 1996; Martin et al., 1990).
By carefully selecting the values of implantation energy and corresponding doses of
carbon ions, a rectangular carbon concentration profile can be achieved for the buried β-
SiC layer using this approach.

Implantation of carbon ions at temperatures of silicon substrate well above room
temperature by in-situ annealing helps to form SiC crystalline layers immediately during
implantation or after annealing at lower temperatures (Durupt et al., 1980; Frangis et al.,
1997; Edelman et al., 1976; Simon et al., 1996; Preckwinkel et al., 1996). High temperature of
the substrate can be achieved also by using beams with high current density of carbon ions
(Reeson et al., 1990; Alexandrov et al., 1986). Treatment of carbon implanted silicon layers
by power ion (Liangdeng et al., 2008; Bayazitov et al., 2003), electron (Theodossiu et al.,
1999) or laser (Bayazitov et al., 2003a, 2003b) beams like thermal annealing also leads to the
formation of a polycrystalline β-SiC layer.
In this paper, the composition and structure of homogeneous SiC
1.4
, SiC
0.95
, SiC
0.7
, SiC
0.4
,
SiC
0.12
and SiC
0.03
layers, received by multiple high-dose implantation of carbon ions with
energies of 40, 20, 10, 5 and 3 keV are investigated. The influence of decay of carbon- and
carbon-silicon clusters during thermal annealing or hydrogen glow discharge plasma
processing on the formation of tetrahedral Si–C-bonds and crystallization processes in
silicon layers with high and low concentrations of carbon, is studied.
2. Experimental
Single-crystal (100) silicon wafers of sizes 7×12×0.4 mm
3

with an electrical resistivity 4–5
Ω⋅сm were implanted by
12
C
+
ions with energies of 40, 20, 10, 5, and 3 keV at room
temperature in vacuum reached by fully oil-free pumping. To prevent sample heating, the
ion current density was kept below 4 μA/cm
2
.
The SiC
x
films structure was investigated by X-ray diffraction using a narrow collimated
(0.05×1.5 mm
2
) monochromatic (CuK
α
) X-ray beam directed at an angle of 5º to the sample
surface. The average crystallite size was estimated from the width of X-rays lines by Jones
method. The surface of the layers was analyzed by Atomic force microscopy (JSPM 5200,
Jeol, Japan) using AFM AC technique. The investigation of morphology and structure of the
SiC
x
layers was carried out by transmission electron microscopy (JEM-100CX, JEOL, Japan).
The IR transmission spectra were recorded in differential regime on double-beam infrared
spectrometer UR-20 (400−5000 cm
-1
). The spectra both at perpendicular incidence of infrared
rays on the sample surface and at an angle of 73
o

with respect to the normal to the sample
surface were measured. The composition of the layers was examined by Auger electron
spectroscopy. The parameters were as follows: incident electron beam of diameter 1 μm,
energy 10 keV, angle of incidence 45°, diameter of scanning region 300 μm, vacuum 1.33
×10
-8
Pa, angle of Ar
+
beam incidence 45°.

Silicon Carbide – Materials, Processing and Applications in Electronic Devices

72
The glow discharge hydrogen plasma was generated at a pressure of 6.5 Pa with a capacitive
coupled radio frequency (r.f.) power (27.12 MHz) of about 12.5 W. The temperature of
processing did not exceed 100°С and it was measured by thermocouple. The processing time
was 5 min.
3. Results and discussion
3.1 Depth profiles of multiple-energy implanted carbon in Si
To produce a rectangular profile of the distribution of carbon atoms in silicon five
energies and doses have been chosen in such a way (Table 1) that the concentration ratio
of C and Si atoms to a depth (up to ~120 nm) was equal to the values of N
C
/N
Si
= 1.0 , 0.8,
0.5, 0.3, 0.1 and 0.03. The calculated profiles N
C
(Burenkov) are the sums of the
distributions (Figs. 1 and 2), constructed for the chosen values of energies (E) and doses

(D) of carbon ions using the values of the ion mean projective range, R
p
(E), and the root-
mean-square deviation, ΔR
p
(Е), (Table 1) and moment S
k
(E
o
), according to Burenkov et al.
(1985). On these figures are presented the calculated profiles N
C
(Gibbons), which are the
sums of Gaussian distributions constructed by using values of R
p
(Е) and ΔR
p
(Е) according
to Gibbons et al. (1975) (LSS):

2
12 2
22
/
()
() exp[ ]
()
p
pp
xR

D
Nx
RR
π
−
=−
ΔΔ
, (1)
where х is the distance
from the surface.
Figs. 1 and 2 also show the experimental profiles N
C
(20°С), N
C
(1250°С) and N
О
(1250°С),
obtained by Auger electron spectroscopy, which show the concentration ratio of carbon and
oxygen atoms (N
C
/N
Si
and N
О
/N
Si
) over the sample depth after implantation and annealing
at 1250°C for 30 min in an argon atmosphere with low oxygen content. Fig. 1 shows that the
average concentrations of carbon and oxygen were: a)N
C

/N
Si
= 1.4 and N
О
/N
Si
≈ 2.6; b)
N
C
/N
Si
= 0.95 and N
О
/N
Si
≈ 2.33; c) N
C
/N
Si
= 0.7 and N
О
/N
Si
≈ 3.0; and as follows from Fig.
2: a)N
C
/N
Si
= 0.4; b)N
C

/N
Si
= 0.12.
Thus, the average carbon concentration over the depth exceeded the corresponding
calculated values (N
C
/N
Si
= 0.1; 0.3; 0.5; 0.8 and 1.0) and led to the formation of layers
SiC
0.12
, SiC
0.4
, SiC
0.7
, SiC
0.95
and SiC
1.4
(Table 2). The average value of the ratio of N
О
/N
Si

exceeded the stoichiometric value for SiO
2
(N
О
/N
Si

= 2), indicating on the saturation of
the surface layer by oxygen. In high-temperature annealing a desorption of carbon atoms
from the surface and an adsorption of oxygen atoms are occurred. There are clear
interfaces SiO
3.0
:SiC
0.7
, SiO
2.33
:SiC
0.95
and SiO
2.6
:SiC
1.4
suggesting the interconnectedness
of these processes. The mechanism of desorption of carbon may be associated with a
diffusion process along the grain boundaries of counter flow of oxygen molecules O
2
into
deep layer and molecules of CO and CO
2
from the layer toward the surface.
Penetration of oxygen deep into the implanted layer until the substrate was also shown in
(Chen et al., 2003). The mechanism of instability of silicon carbide films during high-
temperature annealing in the presence of oxygen is of special interest (Singh et al., 2002).

Process can occur in accordance with the expression
950 1700
22

23 2 2
C
SiC O SiO CO
−°
+ ⎯⎯⎯⎯⎯→+.

The Formation of Silicon Carbide in the SiC
x
Layers
(x = 0.03–1.4) Formed by Multiple Implantation of C Ions in Si

73





0.0
0.5
1.0
1.5
2.0
2.5
3.0
0 40 80 120 160 200
х, nm
N
C
/ N
S i

( N
O
/ N
S i
)
N
С
(Gibbons)
N
C
(40 keV)
N
C
(20 keV)
N
C
(10 keV)
N
C
(3 keV)
N
C
(Burenkov)
N
C
(1250°C)
N
O
(1250°C)
N

С
(5 keV)
N
C
(20°C)
c)


0.0
0.5
1.0
1.5
2.0
2.5
0 40 80 120 160 200
х, nm
N
C
/ N
S i
( N
O
/ N
S i
)
N
C
(Gibbons)
N
C

(40 keV)
N
C
(20 keV)
N
C
(10 keV)
N
C
(3 keV)
N
C
(Burenkov)
N
C
(1250°C)
N
O
(1250°C)
N
C
(5 keV)
N
C
(20°C)
b)


0.0
0.5

1.0
1.5
2.0
2.5
0 40 80 120 160 200
х, nm
N
C
/ N
S i
( N
O
/ N
S i
)
N
C
(Gibbons)
N
C
(40 keV)
N
C
(20 keV)
N
C
(10 keV)
N
C
(3 keV )

N
C
(Burenkov)
N
C
(1250
o
C)
N
O
(1250
o
C)
N
C
(5 keV )
a
)


Fig. 1.
12
С distribution profiles in Si produced by ion implantation (see table 1). (a) SiC
1.4
; (b)
SiC
0.95;
(c) SiC
0.7
. N

C
(Burenkov) and N
C
(Gibbons) are the profiles calculated according to
Burenkov et al. (1985) and Gibbons et al. (1975), respectively, where N
C
(Gibbons)

= N
C
(40
keV) + N
C
(20 keV) + N
C
(10 keV) + N
C
(5 keV) + N
C
(3 keV). N
С
(20
0
С), N
С
(1250
0
С) and
N
О

(1250
0
С) are the Auger profiles of carbon and oxygen, respectively, in a layer after high-
dose implantation and annealing at T = 1250°C for 30 min.

Silicon Carbide – Materials, Processing and Applications in Electronic Devices

74

0.0
0.1
0.2
0.3
0 40 80 120 160 200
х, nm
N
C
/N
Si
b
)
N
C
(Burenkov) .
N
С
(Gibbons)
N
C
(20°C)

N
C
(40 keV)
N
C
(20 keV)
N
C
(10 keV)
N
С
(5 кэВ)
N
C
(3 keV)

0.0
0.1
0.2
0.3
0.4
0.5
0.6
0 40 80 120 160 200
х
, nm
N
C
/N
Si

(N
O
/N
Si
)
N
C
(Burenkov)
N
С
(Gibbons)
N
C
(20°C)
N
C
(40 keV)
N
C
(20 keV)
N
C
(10 keV)
N
С
(5 keV)
N
C
(3 keV)
N

O
(20°C)
а)

Fig. 2.
12
С distribution profiles in Si produced by ion implantation (see table 1). (a) SiC
0.4
; (b)
SiC
0.12
. N
C
(Burenkov) and N
C
(Gibbons) are the profiles calculated according to Burenkov et
al. (1985) and Gibbons et al. (1975), respectively, where N
C
(Gibbons)

= N
C
(40 keV) + N
C
(20
keV) + N
C
(10 keV) + N
C
(5 keV) + N

C
(3 keV). N
С
(20
0
С) and N
О
(20
0
С) are the Auger profiles
of carbon and oxygen, respectively, in a layer after high-dose implantation.

Е, keV 40 20 10 5 3
D
(
SiC
1.0
)
, 10
17
c
m
-2
5.60 1.92 0.990 0.330 0.230
D
(
SiC
0.8
)
, 10

17
c
m
-2
4.48 1.54 0.792 0.264 0.184
D
(
SiC
0.5
)
, 10
17
c
m
-2
2.80 0.96 0.495 0.165 0.115
D
(
SiC
0.3
)
, 10
17
c
m
-2
1.68 0.576 0.297 0.099 0.069
D
(
SiC

0.1
)
, 10
17
c
m
-2
0.56 0.192 0.099 0.033 0.023
D
(
SiC
0.03
)
, 10
17
c
m
-2
0.168 0.058 0.030 0.010 0.007
N
C
(Burenkov) profile
(Burenkov et al., 1985)
R
p
(Е).
nm
120.4 60.0 30.3 16.1 10.5
ΔR
p

(Е).
nm
46.0 28.3 16.9 10.2 7.2
N
C
(Gibbons) profile
(Gibbons et al., 1975)
R
p
(Е).
nm
93.0 47.0 24.0 12.3 7.5
ΔR
p
(Е).
nm
34.0 21.0 13.0 7.00 4.3
Table 1. Values of energy, E, dose, D, projected range, Rp(E), and straggling, ΔRp(E), for
12
C
+

ions in Si, used for constructing a rectangular distribution profile
The Formation of Silicon Carbide in the SiC
x
Layers
(x = 0.03–1.4) Formed by Multiple Implantation of C Ions in Si

75
In (Mandal et al., 2001) reported 100% conversion of SiC into SiO

2
in accordance with the
expression of SiC + O
2
= SiO
2
+ CO/CO
2
.
In the case of high concentrations of carbon (SiC
1.4
, SiC
0.95
, SiC
0.7
and SiC
0.4
) the
experimentally obtained profiles of the carbon atoms were almost rectangular (Figs. 1 and 2)
and close in concentration value to the calculated concentration profile according Gibbons et
al. (1975), while for layers with low carbon concentration SiC
0.12
is closer to the calculated
profile according Burenkov et al. (1985) (Fig. 2). Factors contributing to the excess
concentration of carbon in the surface silicon layer in comparison with the calculated values
by (Burenkov et al., 1985) at high-dose
12
С ion implantation can be considered sputtering the
surface in conjunction with a decrease in the values of Rp due to the presence of strong Si-C-
and C-clusters, double and triple Si–C- and C–C-bonds.

Depth range, nm 40–110 20–120 25–120 10–120 20–120
SiC
х
(Burenkov) SiC
1.0
SiC
0.8
SiC
0.5
SiC
0.3
SiC
0.1

SiC
х
(Gibbons) SiC
1.38
SiC
1.06
SiC
0.67
SiC
0.40
SiC
0.13

SiC
х

(AES) SiC
1.4
SiC
0.95
SiC
0.7
SiC
0.4
SiC
0.12

Тable 2. Average values of carbon concentration x = N
C
/N
Si
in the SiC
х
layers calculated
according Burenkov et al. (1985) and Gibbons et al. (1975) and experimentally obtained by
Auger electron spectroscopy
3.2 Investigation of the structure by electron microscopy
By transmission electron microscopy structural features of the SiC
1.4
, SiC
0.95
and

SiC
0.7
films,

formed in the surface layer of single crystal Si wafers using multiple ion implantation, after
annealing at 1200°C for 30 minutes were investigated. Figure 3 schematically shows a
section of the objects. The investigated area can be divided into three sections: section 1
consists of a layer SiC
x
; section 2 includes a [transition layer "Si-SiC
x
" + layer SiC
x
]; section 3
is a three-layer structure [layer-Si + transition layer "Si-SiCx "+ layer SiCx].

5
SiC
x
SiC
x
SiSi 1
22
33
1
4

Fig. 3. Schematic cross section of the sample under study: (
1) SiC
x
regions, (2) regions of an
intermediate Si-SiC
x
layer, (3) double-diffraction regions, (4) a through hole, and (5) the area

to be analyzed by transmission electron microscopy.
The layers SiC
1.4
, SiC
0.95
and SiC
0.7
are solid, homogeneous, fine polycrystalline films (Figs.
4a-c, 5a, b and 6a, b, light areas, respectively). Some electron diffraction patterns contain
superimposed point (
c-Si) and ring (SiC) electron diffraction patterns (Fig.4b, c, 5b and 6b).

Silicon Carbide – Materials, Processing and Applications in Electronic Devices

76
These patterns were recorded from areas 3 where the objects were presented together with
single- and polycrystalline structures (for example Si + SiC
1.4
). In Figs. 4b, 5b and 6b the
microstructure of sections 1 (light area), 2 (intermediate region) and 3 (dark area) are shown.


1 1 2 3
a) b)

Fig. 4. Electron diffraction patterns and microstructure (×50000) of the SiC
1.4
layer after
annealing at temperature 1200
0

С for 30 min: rings – SiC, point reflections – Si, bright regions
– SiC
1.4
layer, dark regions – c-Si. a) SiC
1.4
region; b) SiC
1.4
regions + intermediate Si– SiC
1.4

layer + с-Si: (
1) SiC
1.4
regions, (2) regions of an intermediate Si– SiC
1.4
layer, (3) double-
diffraction regions.


a) b) c)
1
4
3 2
1

Fig. 5. Electron diffraction patterns and microstructure (×50000) of the SiC
0.95
layer after
annealing at temperature 1200
0

С for 30 min: rings – SiC, point reflections – Si, bright regions
– SiC
0.95
layer, dark regions – c-Si. a) SiC
0.95
region; b) SiC
0.95
regions + intermediate Si–
SiC
0.95
layer + с-Si, c) SiC crystallites in dark-field image regime: (1) SiC
0.95
regions, (2)
regions of an intermediate Si–SiC
0.95
layer, (3) double-diffraction regions, (4) SiC crystallites.
The Formation of Silicon Carbide in the SiC
x
Layers
(x = 0.03–1.4) Formed by Multiple Implantation of C Ions in Si

77

a)
200 n
m
b)
200 n
m
c)

1
2
3

Fig. 6. Electron diffraction patterns and microstructure (×50000) of the SiC
0.7
layer after
annealing at temperature 1200
0
С for 30 min: rings – SiC, point reflections – Si, bright regions
– SiC
0.7
layer, dark regions – c-Si. a) SiC
0.7
region; b) SiC
0.7
regions + intermediate Si–SiC
0.7

layer + с-Si, c) SiC crystallites in dark-field image regime: (
1) SiC
0.7
regions, (2) regions of an
intermediate Si–SiC
0.7
layer, (3) double-diffraction regions
The value of D⋅d of device, calculated from the point patterns of the silicon lattice, was equal
D⋅d = (3.53±0.01) nm⋅mm. The diameters D of the most intensive rings of SiC
1.4
pattern were

10.4, 14.0, 22.9 and 26.9 mm, corresponding to the (002) plane of graphite (d = 0.338 nm) and
silicon carbide β-SiC(111), β-SiC(220), β-SiC(311) (Fig. 4), respectively. The absence of some
rings on the patterns is caused by their weak intensity. The diameters of the most intensive
rings of SiC
0.95
were also 14.0, 22.9 and 26.9 mm and correspond to β-SiC(111), β-SiC(220), β-
SiC(311) (Fig. 5). The diameters of rings of SiC
0.7
diffraction pattern were 14.0, 22.9, 26.9, 28.0,
35.3, 39.7 mm (Fig. 6) and agree well with the calculated values of the diameters of 1, 3, 4, 5, 7
and 9 of the silicon carbide rings, corresponding to a system of planes β-SiC(111), β-SiC(220),
β-SiC(311), β-SiC(222), β-SiC(331), β-SiC(422), (d = 0.2518 nm, 0.1542 nm, 0.1311 nm, 0.1259 nm,
0.1000 nm, 0.0890 nm, respectively). Graphite rings in the electron diffraction patterns of SiC
0.7

and SiC
0.95
are not explicitly observed, which may be due to the absence of excess carbon.
In dark-field image regime the individual grains are visible as bright light spots. For a layer
SiC
0.95
(Fig. 5c), as well as for the layer SiC
1.4
, the presence of small and large grains of size
from 10 to 80 nm was observed. The crystallites have various shapes - globular or plate-like.
For a layer SiC
0.7
(Fig. 6c) crystallites were globular, needle or plate-like and had sizes from
10 to 400 nm after annealing at 1200ºC. Small crystallites have a globular shape. These data
differ from the X-ray diffraction data which was 5

−10 nm (Akimchenko et al., 1977, 1980,
Calcagno et al., 1996). It should be noted that by X-ray diffraction the average crystallite size
was determined. In the layer SiC
0.7
due to the high concentration of carbon atoms one can
expect a large number of stable clusters that hinder the crystallization process. In this case a
significant prevalence of nanocrystals of a few nanometers sizes can be expected that are
making a major contribution to the value of the average grain size. Small nanocrystals
should give a reflection of the size in hundredths and tenths of a millimeter on the pattern, it
is difficult to distinguish them and they are observed as a bright diffuse background
between the larger crystallites (Fig. 6c).

Silicon Carbide – Materials, Processing and Applications in Electronic Devices

78
Transition layer with a lower concentration of carbon between the Si substrate and SiC
0.7
is
not uniform. It can be assumed that excess silicon atoms that are between the large SiC
grains, in the process of recrystallization are combined with the substrate, forming a
sawtooth SiC–Si structure (Fig. 6b).
3.3 Investigation of the structure by X-ray diffraction
After annealing of the SiC
1.4
layer in vacuum 10
-4
Pa at temperatures 1200 and 1400°С in the X-
ray diffraction patterns the intensive lines of polycrystalline silicon carbide are observed (Fig.
7a, b). There are also weak lines of polycrystalline silicon, apparently, from the transition layer
“Si–SiC

1.4
”, where presents excess silicon, forming in the annealing process the Si crystallites.
The integrated intensity of SiC lines after annealing at 1400°C was significantly lower than
after annealing at 1200°C. This may occur because of the disintegration of silicon carbide and
desorption of carbon from the layer at temperatures above 950°С (see section 3.1). The
decrease in phase volume of nanocrystalline SiC at temperatures above 1200ºC can also occur
due to evaporation of silicon in a vacuum, the melting point (1423ºC) of which is less than the
sublimation temperature of carbon (Semenov et al., 2009).


0,4
0,6
0,8
1,0
1,2
1,4
10 30 50 70 90 110
2
θ
, degree
Si (111)
SiС (111)
SiС (220)
2θ, degree
2θ,
degree

c)
2θ, degree
a)

b)
0,2
0,4
0,6
0,8
1,0
1,2
10 30 50 70 90 110
2
θ
, de
g
ree
Si
(
111
)
Si (220
)
SiС (111
)
SiС (200
)
SiС (220
)
SiС (311
)
0,6
0,8
1,0

1,2
1,4
1,6
10 30 50 70 90 110
2
θ
, degree
I, arb. units
Si
(
111
)
Si (220
)
SiС (111
)
SiС (200
)
SiС (220
)
SiС (311
)
SiС (331
)
SiС (422
)

Fig. 7. X-ray diffraction patterns of the SiC
1.4
layer annealed for 30 min at (a) 1200 ºС, (b) 1400°C,

(c) after annealing at 1400°C and processing by glow discharge hydrogen plasma for 5 min.
An absence of SiC crystallites and the corresponding X-ray lines at high annealing
temperatures 900–1100°C indicates a low ability of atoms to diffusion in the layer SiC
1.4
. This
may be caused by high concentrations of stable double and triple bonds like Si=C, Si≡C,
C=C, C≡С, Si=Si, Si≡Si and strong clusters that prevent the diffusion of atoms in the layer
and are decomposed at temperatures of 1200°C and above. An attempt was made to use
processing by the hydrogen glow discharge plasma for the modification of the structural
properties of the SiC
х
films. For this purpose, the SiC
1.4
layer annealed at 1400ºC was treated
by the hydrogen glow discharge plasma (27.12 MHz, 12.5 W, 6.5 Pa, 100°С, 5 min).
After treatment by H-plasma is not taken place the complete destruction of the SiC crystallites
(Fig. 7c). However, the intensity of X-ray lines of SiC decreased in comparison with the
corresponding lines on the diffraction pattern before treatment (Fig. 7b). This demonstrates the
The Formation of Silicon Carbide in the SiC
x
Layers
(x = 0.03–1.4) Formed by Multiple Implantation of C Ions in Si

79
destructive effect of plasma on the structure of the β-SiC crystallites, although the capacity of
plasma was only 12,5 W, and SiC on the hardness scale, according to Knoop holds a high place
(SiC, 2480) after diamond (C, 7000), boron carbide (B4C , 2750), and aluminum boride (AlB,
2500). In addition, the processing by H-plasma for 5 minutes led to complete disappearance of
the lines Si(111) and Si(220), reflected by the crystallites of silicon in the transition layer “film-
substrate" ("SiC-Si"). It follows that the effect of hydrogen plasma propagates on the entire

depth of the implanted layer (200 nm) and involves a intensive penetration of hydrogen ions
through the solid film. Thus, the characteristics of the hydrogen glow discharge plasma and
treatment (27.12 MHz, 12.5 W, 6.5 Pa, 100°С, 5 min) were sufficient to break the Si-Si and Si-C
tetrahedral bonds and can be used to decay the stable carbon and carbon-silicon clusters.
Immediately after implantation and annealing of the SiC
0.95
layer at 1100ºC and below (Fig.
8a) on the diffraction patterns not seen any lines of polycrystalline phases. A weak line of β-
SiC(111) after annealing at 1150ºC is observed, which becomes more pronounced after
annealing at 1250
0
С (Fig. 8b). In contrast to the layers SiC
1.4
and SiC
0.95
, the diffraction
pattern of the SiC
0.7
layer (Fig. 8c) demonstrates the presence of β-SiC(111) line and weak Si
(111) line after annealing at relatively low temperature of 1000ºC for 30 minutes, due to the
decrease of carbon concentration and, consequently, the concentration of stable carbon
clusters. Annealing at 1100ºC led to an increase in the intensity of the β-SiC(111) and Si(111)
lines (Fig. 8d) and to an appearance of β-SiC(200), β-SiC(220), β-SiC(311) и Si(220) lines,
which indicates an improvement in the crystallite structure. Line intensity from the SiC
0.7
layer shows a preferred content of polycrystalline β-SiC phase after annealing in comparison
with the phase content of poly-Si, which is caused by a relatively high concentration of
carbon. The average crystallite size of β-SiC and Si was about 3
–7 nm in the planes:
β-SiC(111) ~ 3 nm, β-SiC(220) ~ 6.5 nm, Si(111) ~ 4.5 nm.

d)
c)
0
0,1
0,2
0,3
0,4
0,5
0,6
10 30 50 70
2
Η
, degree
I
, ar
b
.un
.
Si (111)
Si (220)
SiС (111)
SiС (200)
SiС (220)
SiС (311)
b)
а)
0,6
0,7

0,8
0,9
1,0
10 20 30
2

, degree
I, arb.un
.
SiС (111
)
Si (substrate
)
θ, degree 2θ, degree

Fig. 8. X-ray diffraction patterns of the SiC
0.95
layer annealed at (a) 1100ºС, (b) 1250°C and of
the SiC
0.7
layer annealed at (c) 1000ºС, (d) 1100°C , for 30 min obtained using X-ray chambers
RKD (a, b) and RKU-114M1 (c, d). Intensity curves correspond to X-ray patterns (b, d).

Silicon Carbide – Materials, Processing and Applications in Electronic Devices

80
In the layer SiC
0.4
the presence of polycrystalline phases of SiC and Si after implantation and
annealing at 1000, 1100 and 1250°C are revealed. High intensity peaks of SiC after annealing

at 1100ºC (Fig. 9a) indicates on intensive process of crystallization of β-SiC due to lower
content of stable clusters in comparison with the layers SiC
1.4
, SiC
0.95
and SiC
0.7
. High
amplitude of β-SiC peaks after annealing at 1250
о
С confirms this assertion. In diffraction
pattern of SiC
0.12
layer after implantation a broad diffuse line of amorphous silicon Si(111)
(θ = 14.3º) is observed (Fig.10a). After annealing at 800ºC a narrowing of this line, at 900ºC -
a sharp narrowing of the line Si(111) and the appearance of Si(220) and Si(311) lines of
polycrystalline Si phase, as well as two weak lines of β-SiC, at 1250ºC - a decrease of the
integrated intensity of Si lines, are observed.

0,4
0,6
0,8
1,0
1,2
1,4
1,6
10 20 30 40
θ
,
de

g
ree
I, arb.un
.
SiС (111
)
Si (substrate
)
Si (111
)
Si (220
)
Si (311
)
SiС (220
)
SiС (311
)
b)
0,1
0,3
0,5
0,7
0,9
1,1
1,3
515253545
θ
, de
g

ree
I, arb.un
.
Si (111
)
SiC (111
)
Si (220
)
SiС (220
)
Si (311
)
а)

Fig. 9. X-ray diffraction patterns of the SiC
0.4
layer annealed for 30 min at (a) 1100ºС and (b)
1250°C.
The implantation of carbon causes the formation of weakly ordered set of randomly
oriented Si regions of size ~1.5 nm. Temperature dependence of the average crystallite size
Si (Fig. 11a, curve 1) and SiC (curve 2) in plane (111) shows that at low temperatures the
curve 1 is characterized by slow growth in the size of weakly ordered Si nanocrystals. Their
transformation into well ordered Si crystallites at 800ºC is taken place. Average size of Si
crystallites is 47 nm after annealing at 1250ºC. The formation of β-SiC crystallites prevent a
complete recrystallization of layer at this temperature.
The integrated intensity curve of Si(111) line, which is proportional to the phase volume,
has three sections in the temperature ranges of 20– 800ºC, 800−900ºC and 900−1250ºC (Fig.
11b). In the range of 20-800ºC I
int

decreases, indicating a decrease in the total volume of
nanocrystals Si, although their sizes increase up to 2.2 nm. The growth of weak ordered Si
crystallites is accompanied by the displacement of carbon atoms in the surrounding space,
which no longer contribute to the intensity of the Si (111) line due to increased
concentration of carbon. In addition, there is recrystallization of silicon near the substrate.
It is assumed that, after annealing at 800ºC (Figs. 11b and 12) approximately
k
800
≈ [I
int
(800ºС)/I
int
(20ºС)]×100% = 72% of silicon atoms in the layer is incorporated into
crystallites of Si. The remaining atoms of Si (28%) are in an amorphous mixture of Si and
C atoms, or reunited with the substrate.
The Formation of Silicon Carbide in the SiC
x
Layers
(x = 0.03–1.4) Formed by Multiple Implantation of C Ions in Si

81

0,2
0,7
1,2
1,7
2,2
2,7
10 20 30 40 50 60
θ

, de
g
ree
Si (111)
Si (220)
Si (311)
Si
(
422
)
Si
(
511
)
Si
(
331
)
Si (531)
Si
(
matrix
)
SiC (111)
SiC (311)
SiC (220)
SiC (422)
SiC (331)
SiC (400)
0,0

0,5
1,0
1,5
2,0
2,5
10 20 30 40 50 60
θ
,
de
g
ree
Si (111)
Si (220)
Si (311)
Si (422)
Si (511)
Si (531)
Si (matrix)
SiC (111)
SiC (311)
1,1
1,6
2,1
2,6
3,1
3,6
10 20 30 40 50 60
θ
,
degree

I, arb. units
Si (111)
Si (matrix)
a) b) c)
θ, degree
θ, degree
θ, degree

Fig. 10. X-ray diffraction patterns of the SiC
0.12
layer (a) before and after annealing for 30 min
at (b) 1100ºС and (c) 1250°C.

0,00
0,04
0,08
0,12
0,16
20 300 600 900 1200
Temperature,
o
С
I
int
, arb.un.
2
3
1
b)
0

10
20
30
40
50
20 300 600 900 1200
Temperature,
o
С
Crystallite sizes, nm
.
2
3
1
a)

Fig. 11. Average sizes (a) of Si and SiC crystallites in the (111) plane and (b) the integrated
intensities of the Si(111) and SiC(111) lines after implantation of carbon ions in silicon and
annealing. 1 − Si (for layer SiC
0.03
), 2 − Si (for layer SiC
0.12
), 3 − SiС (for layer SiC
0.12
).

Silicon Carbide – Materials, Processing and Applications in Electronic Devices

82
Fig. 12 schematically shows the variation of the structure of the layer, its phase composition

and phase volume, as well as the average grain size as a function of annealing temperature
(a) and scheme of Si and SiC crystallite formation in this layer (b). The diagram is based on
the curves in Fig. 11a, b for SiC
0.12
. Regions "amorphous Si-C mixture” and "c-Si" are formed
adhering to the relation: "amorphous Si-C mixture” + "c-Si" = 100% − (“poly-SiC” + “poly-
Si"+"amorphous -Si").


Fig. 12. Crystallization of the SiC
0.12
layer: (a) the phase volumes at various annealing
temperatures and (b) the formation of crystalline Si and SiC crystallites in the temperature
range 900−1000°C.
In the range of 800−900°C the increase in both size crystallites Si (from 2.2 to 4.7 nm) and the
volume of polycrystalline phase of Si (Fig. 11a and 12) is observed. This is due to the
formation of SiC crystallites at 900°C in the regions of carbon accumulation and the joining
of excess silicon atoms to Si crystallites. So, k is increased: k
900
≈ 82%. This leads to a decrease
in Si−C-mixture volume (Fig. 12). After annealing at 1000ºC and 1250ºC, an increase in the
phase volume and crystallite sizes of β-SiC as well as a decrease in k due to recrystallization
of regions near the Si substrate, are taken place: k
1000
≈ 73% and k
1250
≈ 46%. Si crystallite
sizes increased almost 10 times (up to 47 nm). Probably, there is destruction of defective
silicon crystallites and the uniting of their atoms into crystallites with a perfect structure or
with the substrate. As can be seen from the diagram, after annealing at 1200°C in the layer

SiC
0.12
~50% of silicon atoms are incorporated into Si crystallites with an average size of 25
nm, 25% of Si atoms are included into β-SiC crystallite with size of 5 nm and 25% are joined
with Si substrate. Since the unit cell volume for Si is twice more than the cell volume for SiC,
the volume of the surface layer of Si after implantation and annealing of carbon should not
change appreciably, and the above relations can be regarded as a volume ratio of phases.
During annealing at 1400ºC there is the destruction of most of the silicon crystallites and
connection of their atoms to the substrate. The intensity of the lines Si was significantly
lower than the intensities of β-SiC lines (Fig. 13). It is assumed that recrystallized Si layer
with ingrained in him crystallites of β-SiC and Si on a silicon substrate was obtained.
During a growth of the crystallite size is manifested basic thermodynamic law: in an isolated
system, the processes occurring with increasing free energy, is prohibited. Combining the
The Formation of Silicon Carbide in the SiC
x
Layers
(x = 0.03–1.4) Formed by Multiple Implantation of C Ions in Si

83
two grains is taken place, if it is accompanied by a gain in energy which is greater than its
costs for the destruction of the crystallites. In this case, the redistribution of atoms with a
change in chemical bond lengths and angles between them is taken place, in order to select a
more favorable energy state, which is the states with tetrahedral oriented bonds,
characteristic of crystalline phases of Si and SiC.

0,25
0,35
0,45
0,55
0,65

0,75
10 30 50 70 90
2
θ
, degree
I, arb.units
Si (111)
.
Si (220)
.
SiС (111)
.
SiС (200)
.
SiС (220)
.
SiС (311)
.

Fig. 13. X-ray diffraction pattern of the SiC
0.12
layer after annealing at 1400°C for 30 min.
The X-ray diffraction results are in accordance with the data of Auger electron spectroscopy,
whereby the concentration of carbon atoms in a layer is N
C
/N
Si
= 0.12/1 (Fig. 2). Then the
maximum possible ratio of atoms, forming part of SiC and Si, will be: N
SiC

/N
Si
= 0.24/0.88 =
0.27. As seen in Figure 11b, the maximum quantity of poly-SiC, obtained at 1250ºC, was
I
int
(SiC) = 0.040, and poly-Si was I
int
(Si)

= 0.131 at 900ºС, i.e., I
int
(SiC)/I
int
(Si)

= 0.040/0.131 =
0.30, which is comparable with the data of Auger electron spectroscopy.
Similar features are observed for a layer with a lower carbon concentration SiC
0.03
.
Immediately after the implantation the same broad diffuse line of amorphous silicon Si(111)
at θ = 14.3º is observed (Fig. 14). Increase of annealing temperature above 800ºC causes
narrowing of this line, increasing the number and amplitude of the line intensity which
reaches its maximum at 1100ºC. However, in the X-ray diffraction patterns of SiC
0.03
layer no
lines of polycrystalline silicon carbide in the whole temperature range are observed due to
insufficient concentration of the carbon atoms to form a large number of SiC crystallites.
Annealing at temperatures above 1100º C reduces the intensity of Si lines and results their

disappearance after annealing at 1250ºC due to recrystallization of the layer.
For comparison, in Table 3 the grain sizes of silicon and silicon carbide in plane (111) in
layers Si(111) and β-SiС(111) are given. Dimensions of weakly ordered regions in SiC
0.03

layer, contributing to the intensity of the Si(111) line, exceed the same values for SiC
0.12
. This
is associated with a lower concentration of carbon and larger volume of Si regions with
extremely low concentrations of carbon. At temperatures of 1200 and 1250ºC a decrease in
the average crystallite size up to 10 nm is observed, which is associated with the process of
recrystallization near the substrate.